Regular Article
Alloying Effects on the Microstructure of Fe-1mass%M Binary Alloys Treated by Austenitic Nitriding and Quenching Treatment
2024 年 64 巻 2 号 p. 226-234
詳細
2024 年 64 巻 2 号 p. 226-234
Surface microstructures were investigated in pure iron and Fe-1mass%M (M = Mn, Cr, Al, Si) alloys gaseous-nitrided at 1123 K and quenching to reveal the alloying effects on surface hardening by nitrogen (N) martensite. Thicker hardened layers with higher hardness than pure iron were obtained in the Mn-added alloys whereas the additions of Si and Al lead to increase the surface hardness with reduction of the hardened layer thickness. On the other hand, adding Cr decreases both the hardness and thickness of the hardened layer. No precipitation of alloy nitride is observed in austenite nor internal ferrite region in Mn-added alloy. Meanwhile, CrN(B1) and AlN(wurtzite) particles are dispersed in ferrite and austenite regions in the Cr- and Al-added alloys, respectively. Unlike those alloys, (austenite+ α-Si3N4) lamellar structure is formed in the Si-added alloy followed by martensite transformation of the high temperature austenite during quenching. Phase diagrams of Fe–M–N systems can consistently describe those alloying effects on the microstructure evolution.
To improve fatigue strength and wear resistance of steel parts, carburizing and nitriding have been widely used in industry for decades; the former is surface hardening treatment by high carbon (C) martensite, while the latter hardens the surface by precipitation of fine alloy nitrides or cluster and the compound layer. Despite many years of research, the weaknesses of carburizing and nitriding remain: carburizing has significant heat treatment distortion due to high heating temperature and martensite transformation during quenching, and nitriding needs a long heat treatment time to obtain appropriate thickness of surface-hardened layer. Therefore, austenitic nitriding,1) known as nitriding and quenching (N-Q) treatment2) has been developed to improve those disadvantages.
Austenitic nitriding is performed above the eutectoid temperature of the Fe–N system (863 K) to form high N austenite near the surface, followed by quenching to martensite transformation. The austenitic nitriding process (1) has smaller heat treatment distortion than carburizing because the heating temperature is frequently lower than Ae1 temperature corresponding to a carbon-contained steel in the internal region, and austenite transformation does not occur,3) (2) has a relatively shorter treatment time than nitriding treatment (thicker thickness of the hardened layer) because the temperature is higher than nitriding treatment1,4) and (3) hardens the surface by formation of martensite. So alloying elements to form alloy nitride, such as V, Ti, Cr, etc., are not necessary in austenitic nitriding unlike to nitriding treatment of ferritic/martensitic alloys.1,2,3,5) Although austenitic nitriding has been reported to give superior wear and corrosion resistance to carburizing3) and higher rotating bending fatigue strength than nitrocarburizing,6) a deeper understanding of microstructure evolution during the treatment is needed for further alloy/process development.
Austenite growth kinetics and hardness of N martensite after quenching have been investigated using pure iron.2,7) Austenite growth kinetics becomes faster at higher austenitic nitriding temperature or higher nitriding potential and follows parabolic law when the formation of the void at the surface is avoided.2,7) The hardness of martensite in Fe–N alloys increases at higher N content and decreases above 1.2 mass%N due to retained austenite e.2,8) This composition dependency is similar to C martensite,9) but N martensite is softer than C martensite at the same composition.2,8)
On the other hand, alloying elements should influence austenite stability and hence austenite growth kinetics. In addition, it can induce alloy nitride precipitation during austenitic nitriding. Pavlick et al.10) investigate the effects of V addition on austenite growth kinetics. They found that V addition promotes austenite penetration along the ferrite grain boundary near the surface while retarding the thickening kinetics of an austenite layer from the surface. However, there is no systematic investigation on alloying effects on surface hardening and nitride precipitation in austenitic nitriding. Therefore, this study investigates the microstructure and hardness of Fe-1mass%M binary alloys after austenitic nitriding to clarify the influence of alloying elements on microstructure evolution.
Pure iron and Fe–M binary alloys, which contain 1 mass% of Mn, Cr, Si, or Al, were used in this study. The alloy compositions are shown in Table 1. After a homogenization treatment at 1423 K for 86.4 ks, plate specimens with the dimensions of 15 × 8 × 3 mm3 were cut, mechanically polished, and cleaned in acetone before nitriding.
| C | Si | Mn | Cr | Mo | Al | |
|---|---|---|---|---|---|---|
| Pure iron | 0.011 | 0.02 | 0.02 | 0.01 | – | – |
| Fe-1Mn | 0.001 | 0.01 | 0.98 | – | – | – |
| Fe-1Cr | 0.010 | 0.02 | 0.02 | 0.99 | 0.01 | – |
| Fe-1Si | 0.001 | 0.95 | 0.01 | – | – | – |
| Fe-1Al | 0.002 | – | – | – | – | 0.99 |
The specimens were gaseous-nitrided in a vertical quartz tube furnace with an inner diameter of 22 mm using H2/NH3 mixture at 1123 K. The total gas flow, the total pressure of the mixture gas, and the partial pressure of NH3 were 100 ml/min, 1 atm, and 0.1 atm, respectively, corresponding to nitriding potential, rN = 0.117/atm−0.5. At this rN, iron nitrides such as ε-Fe2–3N and γ’-Fe4N are not formed on the surface of specimens. In addition, N activity at this nitriding potential estimated from the equilibrium constant of NH3 decomposition1) greatly exceeds unity, suggesting that the formation of the void could occur by the precipitation of N2 gas. However, microstructure observation confirms no void formation in the nitrided region due to insufficient time for this reaction. After the nitriding treatment for various periods, the samples were quenched into the water. For the subsequent investigation, nitrided specimens were cut, and their cross-sections were mechanically polished for microstructure observation and hardness and N content measurements.
Microstructure observations were made by optical microscopy (OM), scanning electron microscopy (SEM; JEOL JSM-7001 operated at 5 kV), and transmission electron microscopy (TEM; FEI CM-300 operated at 300 kV, JEOL JEM-2100Plus operated at 200 kV). The prepared specimens were etched with 3% nital for OM and SEM observations. Hardness profiles on cross-sections were measured by a micro-Vickers hardness tester with a load of 100 g. To compare the effects of compositions on the hardness between N and C martensite, five Fe–C binary alloys with C contents of 0.35 mass%, 0.60 mass%, 0.75 mass%, 0.90 mass%, and 1.05 mass% were also used in this study. Those Fe–C alloys were austenitized at 1273 K for 1.8 ks, followed by water quenching to measure the hardness of Fe–C martensite.
N content profiles on cross-sections were determined using a field emission electron probe microanalyzer (FE-EPMA; JEOL JXA-8530F operated at an accelerating voltage of 15 kV and current of 100 nA). A three-dimensional atom probe (3DAP) was used to analyze N atoms in solution and nano-scale alloy nitrides. Specimens for TEM observations and 3DAP analysis were made using a focused ion beam with a microsampling method (FIB; FEI Quanta 3D). 3DAP analysis was carried out using the LEAP 4000HR produced by CAMECA. The specimen temperature was 50 K, and the pulse fraction was 20% for 3DAP analysis. The peaks at mass to charge ratio (m/n) = 7, 14, and 15 were assigned to N ions (N2+, N+), m/n = 18.6, 27 and 28–29 to (Fe3+, Fe2+), m/n = 34–36 to (FeN)2+, m/n = 17.3 and 25–26.5 to (Cr3+, Cr2+), m/n=32–33.5 to (CrN)2+, m/n = 9 and 13.5 to (Al3+, Al2+), m/n = 27.5 to (Mn2+). In the case of Fe-1Si alloy, peaks of 28Si2+ and 14N+ ions are overlapped at m/n = 14, and those peaks of the mass spectrum were decomposed into respective ions according to the natural isotope abundance. On the other hand, the overlap of N+ and
Figure 1(a) shows a cross-sectional OM image of the pure iron nitrided at 1123 K for 36 ks and subsequently quenched. A darkly etched layer near the surface was austenite at nitriding temperature and finally transformed into martensite after quenching, which will be referred to nitrided zone (γ). The white area below the nitrided zone is the original ferrite structure. After the nitriding for 36 ks, the austenite phase grows to a thickness of about 450 μm. Figure 1(b) shows the corresponding hardness- and N content-depth profiles. The hardness of the nitrided zone reaches approximately 700 HV near the surface and gradually decreases with depth, then drops sharply at the growth front of the nitrided zone. The N content is 0.46 mass% near the surface and decreases with depth, similar to the hardness profile.
Figure 2(a) shows hardness-depth profiles of Fe-1M alloys compared with that for pure iron. A short vertical bar on each hardness profile indicates the position of the interface between the nitrided zone (γ) and ferrite. Analogous to pure iron, pronounced hardening occurs in all specimens’ nitrided zone. The nitrided zones of the Mn-added alloys have higher hardness and larger thickness than those for pure iron. On the other hand, Fe-1Al and Fe-1Si show a thinner but harder nitrided zone than those for pure iron. The Al addition, notably, leads to significant hardening of more than 800 HV. The Cr addition decreases both hardness and thickness. Thus, it is concluded that adding Mn accelerates the growth of austenite, while Si, Cr, and Al additions retard it. Figure 2(b) shows the corresponding N content-depth profiles. All the alloys have higher N content near the surface than pure iron. In particular, significant N uptake occurs near the surface of the Al and Si-added alloys and ferrite of the Al-added alloy. On the other hand, the increment of N content in the nitrided zone and ferrite is marginal in the Cr and Mn-added alloys.
Figure 3 shows optical micrographs of the region close to the interface between the nitrided zone and ferrite in Mn, Si, Cr, and Al-added alloys. In the Mn-added alloy (Fig. 3(a)), lath martensite structure is observed in a darkly etched area, but no alloy nitride precipitation is observed despite higher N content than that of the pure iron. Dispersion of alloy nitrides is observed both in the nitrided zone and ferrite of Al- and Cr-added alloys (Figs. 3(b) and 3(c)). On the other hand, Si-added alloy contains lamellae structure in the nitrided zone and coarse precipitate in ferrite, as shown in Fig. 3(d).
Those nitrides are characterized using SEM, TEM, and 3DAP, as shown in Figs. 4, 5, 6, 7, 8, 9. Figure 4 shows nitrides formed in the Fe-1Al alloy. In the ferrite region below the nitrided zone, particles lying in rows are observed in Fig. 4(a). The bright-field TEM image and corresponding selected area diffraction analysis indicate that the precipitates are AlN with a wurtzite structure having a Pitsch-Schrader orientation relationship (0001)AlN//(011)α, [1010]AlN//[011]α, [1210]AlN//[100]α to ferrite matrix. SEM and bright-field images in Figs. 4(e) and 4(f) show that the nitrided zone contains many precipitates lying in rows in martensite structure, and the particles are identified as AlN (wurtzite) by TEM analysis. Similar dispersion of AlN particles in martensite to ferrite indicates that austenite should grow by inheriting AlN particles in ferrite.
Figure 5(a) is a bright-field TEM image of a precipitate in a ferrite matrix in front of the nitrided zone of the Fe-1Cr alloy. The SADP analysis in Figs. 5(b) and 5(c) reveals that the disk-shaped precipitate lying on (100) α is CrN with NaCl (B1) structure and has the Baker-Nutting orientation relationship, (001)CrN//(001)α, [100]CrN//[110]α, [010]CrN//[110]α to ferrite matrix.16,17) CrN (B1) precipitation on {001}α is frequently reported in ordinary ferritic nitriding of Cr-added alloys at a temperature below 863 K.16,18,19,20)
Figures 6(a) and 6(b) show a bright-field TEM image of the nitrided zone at a depth of 50 μm in the Fe-1Cr alloy. Arrows indicate that the nitrided zone contains ferrite grains with lower dislocation density than the surrounding martensite. Most ferrite grains contain CrN particles, as shown in the enlarged image and SADP analysis of Figs. 6(b) and 6(c). As will be discussed later, α cannot exist thermodynamically at nitriding temperature. Therefore, these results suggest that intragranular ferrite nucleation during quenching after austenitic nitriding is induced by the CrN particle in γ.
Figure 7(a) shows a bright-field TEM image of precipitates in the ferrite matrix of the Fe-1Si alloy. The precipitates are identified as α-Si3N4 (trigonal, P31c, a=0.7818 nm, c=0.5591 nm21)) by analysis of selected area diffraction pattern as shown in Fig. 7(b). The α-Si3N4 precipitates appear to have polygonal shapes with some developed facets.
Figure 8(a) shows a secondary electron image and elemental mappings of N and Si of the nitrided zone (γ) at a depth of 50 μm in Fe-1Si alloy, taken by FE-EPMA. The enrichment of N and Si at precipitates in the lamellae suggests that those precipitates are Si nitrides. The bright field TEM images of the lamellar region are shown in Figs. 8(b) and 8(c). As in the case of ferrite region, the crystal structure of those precipitates corresponds to α-Si3N4, as seen in Figs. 8(d) and 8(e). It is also found that the counterpart forming the lamellar structure is the martensite structure in Fig. 8(c). Since this martensite is presumably formed by quenching, it is concluded that (austenite + α-Si3N4) lamellar structure evolves in the Fe-1Si alloy during nitriding.
3.4. N Content in the SolutionFigure 9 shows three-dimensional mappings of atoms of N and alloy elements taken from 3DAP of the nitrided zone (γ) at 50 μm-depth from the surface of each specimens. Table 2 shows N and M contents in γ region, excluding any alloy nitrides present. Plate-shaped regions enchied by N are observed, as shown in the pure iron (Fig. 9(a)). N content of these regions are approximately 11 at%, which closely matchs the composition of Fe16N2 (11.1 at%). This indicates the precipitation of Fe16N2 in martesite during quenching by autotempering as is observed in C martensite.22) Autotempering is weakened in the alloyed specimens (Figs. 9(b)–9(e)). Homogeneous Mn distribution and the Mn content nearly equal to the nominal one indicate no Mn nitride precipitation in the Fe-1Mn alloy. 3DAP measurements of the pure iron and Fe-1Mn alloy yield 17% higher N content than FE-EPMA. This discrepancy is considered an artifact of 3DAP measurement as explained in the experimental procedure,12,15) and this factor was multiplied to calibrate the underestimation of N content in 3DAP measurement, and only calibrated N contents are shown in Table 2. N content in the Fe-1Mn and Fe-1Al alloys in γ is slightly higher than pure iron. AlN particles are included in the measured region as shown in Fig. 9(e), and Al content in γ is negilibley small (Table 2). Similarly the lower Si and Cr contents from the nominal ones suggest that some of those elements are consumed to precipitate alloy nitrides, especially for the Al and Si. On the other hand, N in solution increases by the addition of Mn, Al, and Cr but decreases by the Si addition.
| Pure iron | Fe-1Mn | Fe-1Al | Fe-1Cr | Fe-1Si | |
|---|---|---|---|---|---|
| N content | 0.54 | 0.61 | 0.61 | 0.55 | 0.47 |
| M content | – | 1.02 | 0.00 | 0.30 | 0.05 |
Alloying effects on microstructure and hardness distribution in austenitic nitriding and quenching treatment at 1123 K were investigated in this study. Mn, Al, or Si addition increases surface hardness compared with pure iron, while Cr addition reduces surface hardening. Meanwhile, adding all Mn, Al, Cr, and Si increase surface N contents in the nitrided region (γ) by the precipitation of alloy nitrides and the increment of N content dissolved in austenite except Si. Cr, Si, and Al nitride precipitation is observed in the ferrite matrix and austenite regions. In this section, those microstructure evolution will be discussed.
In Fe-1Si alloy, α-Si3N4 is precipitated in ferrite, and the nitrided zone (γ) consists of (γ + α-Si3N4) lamellar structure as shown in Fig. 8. Precipitation of amorphous Si3N4 in ferrite matrix is frequently reported in Si-added alloy nitrided at below 853 K.23,24,25,26,27,28) Amorphous Si3N4 instead of the thermodynamically stable crystalline Si3N4 is formed due to the relatively low energy of the interface between amorphous Si3N4 and ferrite matrix than that between crystalline Si3N4 and ferrite matrix.24,25) It was also pointed out that the volume misfit of α-Si3N4 is extremely large in comparison to other alloys nitrides (51%(CrN), 77%(AlN-wurtzite), 108%(α-Si3N4)1,29)), which causes sluggish precipitation kinetics1. The preferential precipitation of α-Si3N4 over amorphous Si3N4 in the present study suggests that raising nitriding temperature enhances precipitation of α-Si3N4 over the amorphous phase due to easier accommodation of volume misfit by the faster diffusion of vacancy and promoted plastic deformation.
On the other hand, formation of (γ + Si3N4) lamellar in the Fe-1Si alloy is, at least to the authors’ knowledge, the first to be observed in the austenitic nitriding treatment. A similar reaction occurs during carburizing of alloy steels and is known as δ-pearlite transformation forming by eutectoid transformation.30,31,32,33,34) Figure 10(a) shows an isothermal phase diagram of the Fe–Si–N system at 1123 K calculated using Thermocalc with TCFE12 database. N activity with a reference of 1 atm N2 gas is the vertical axis since N activity is a controlling variable in nitriding treatment. 1 mass%Si is hypereutectoid composition, and Si3N4 is precipitated first in the ferrite matrix. Subsequently, the remaining ferrite matrix is transformed into (γ +Si3N4) at higher N activity. This microstructure evolution is in good agreement with the experiment.
In the Cr-added alloy, disk-shaped CrN(B1) is precipitated in the ferrite matrix as frequently observed in nitriding at around 823 to 873 K temperature but much smaller than the present study.19,35) According to the isothermal phase diagram shown in Fig. 10(b), precipitation of CrN proceeds in a ferrite matrix before austenite formation via the peritectoid reaction. It is also seen that α region is reduced by the addition of Cr. Eventually, ferrite disappears in Fe-1Cr alloy at sufficiently high N activity where austenite is stable in pure iron. Therefore, ferrite grains surrounding CrN (B1) shown in Fig. 6 are unlikely to exist in the nitrided zone (γ) during nitriding, and ferrite should be formed during quenching. The 3DAP analysis of ferrite/γ(M) interface in Fig. 11 reveals that Cr is not undergoing partitioning at this interface, despite the expectation that Cr should experience enrichment in the γ phase compared to ferrite in an equilibrium state in ThermoCalc calculation. Consequently, the absence of Cr partitioning at the interface supports that ferrite does not form during nitriding, but rather during the cooling process. Ferrite nucleation at CrN (B1) is probably induced by the relatively lower interface energy between ferrite and CrN (B1) in cooperation with the somewhat higher energy of the incoherent interface between austenite and CrN (B1) inherited from ferrite.
The isothermal phase diagram of the Fe–Al–N system shown in Fig. 10(c) suggests precipitation of AlN in the ferrite matrix, followed by α + AlN→γ + AlN transformation, consistent with the experiment. Precipitation of AlN (wurtzite) in ferrite matrix in addition to metastable AlN(B1) is frequently reported in ferritic nitriding at lower temperature.16,17,36,37,38,39) Figure 10(c) shows that the (ferrite + AlN) phase region extends to lower N activity because AlN has much higher thermodynamic stability than Si3N4 and CrN. Consequently, the nucleation driving force per one mole of substitutional atoms from the ferrite matrix evaluated at the highest N activity in ferrite region is 5900 J/mol for AlN, that is much larger than those for Si3N4 (160 J/mol) and CrN (890 J/mol). In contrast, precipitation of AlN(wurtzite) is less intensive than CrN, as shown in Fig. 3. Slower precipitation kinetics of AlN(wurtzite) than CrN has been explained by a more significant volume misfit of AlN(wurtzite) than CrN(B1).1) On the other hand, recently, we focused on the solute element interaction between N and substitutional in ferrite matrix40,41) and found that solute clustering is induced by attractive interaction in the case of nitriding of Cr, Ti, and V-added alloys. Meanwhile, Al and N have weak or even repulsive interaction in the ferrite matrix, that can also result in the sluggish precipitation kinetics of AlN(wurtzite) irrespective to the significantly large driving force.40,41)
Alloying effects on N content in austenite solution are also important since N content in martensite significantly impacts martensite hardness.2,8) Figure 12(a) shows variations of total N content with the amount of alloying element at 1123 K. Constant N activity of 18 is assumed to reproduce N content (0.52%) in pure iron. Total N content in each of Al-, Si- and Cr-added alloys increases with increasing the amount of element addition due to the precipitation of alloy nitrides. In contrast, Mn–N attractive interaction increases total N content by Mn addition. The increment of total N content by Mn and Cr addition agrees well with EPMA measurement.
On the other hand, N uptake in Al or Si-added alloy exceeds the calculated N contents significantly. In the nitriding of ferritic alloys, N uptake is sometimes higher than the equilibrium N composition contained as alloy nitrides and solution in the ferrite matrix, called excess N.42) However, excess N is not reported in austenitic nitriding. A systematic study will be needed to clarify the abnormally high N uptake in Al- and Si-added alloys.
Figure 12(b) represents variations in N dissolved in austenite and total N content with amount of alloying elements. Cr addition increases but Si addition decreases N content in austenite at small amounts of addition due to attractive and repulsive interaction against N, respectively. N content in austenite becomes constant above the solubility limit for alloy nitrides. The solubility limit of AlN is tiny (less than 1 ppm), and the effect of Al addition on N content austenite is negligibly small. On the other hand, Mn nitride is not formed in the condition calculated, and Mn–N attractive interaction increases N content in austenite, as explained above. This tendency is consistent with the 3DAP measurement shown in Table 2 except Al, where Al increases N content in austenite as much as Mn addition. In ferritic nitriding, the absorption of nitrogen at the ferrite/alloy nitride interface and the enhancement of nitrogen solubility resulting from elastic strain induced by the precipitation of alloy nitrides, have been proposed as factors contributing to the excess of nitrogen.42) The smaller interface area, due to the larger size of alloy nitrides in austenitic nitriding compared to ferritic nitriding, implies that the latter phenomenon might account for the observed increase in nitrogen content in solution due to the addition of Al during austenitic nitriding.
4.2. Alloying Effect on Hardness of MartensiteFigure 13 shows N and C martensite hardness in this and the previous studies as a function of C or N content.2) The hardness of Fe–N martensite increases up to approximately 800 HV with increasing N content, meanwhile it is softer than that in Fe–C alloys. Tsuchiyama et al.8) also reported that N martensite is softer than C martensite. They showed that the volume fraction of retained austenite in the Fe–N system is lower than in the Fe–C system, which is inconsistent with lower hardness, and concluded that a part of the softening is due to low dislocation density in N martensite.8) Lüthi et al.43) performed a DFT calculation on interstitial elements’ interaction with screw dislocation and found that N interaction to dislocation is much weaker than C, which could be another reason for the softness of N martensite.
Solid and open symbols for 1Al, 1Si, and 1Cr-added alloys represent N contents measured using EPMA and 3DAP, respectively. The former indicates total N content, including N atoms in alloy nitrides and N atoms dissolved in austenite, and the latter is only N content dissolved in austenite. The 1Mn-added alloy exhibits nearly the same tendency as the Fe–N alloy. Hence, enhanced hardening by the Mn addition in Fig. 2 should be originated from the higher N content in austenite caused by Mn–N attractive interaction. The hardness of nitrided zone (γ) in the Cr-added alloy is lower than those of pure iron at the same total N content or N content dissolved in austenite due to ferrite formation during quenching, as shown in Fig. 6. Addition of Al and Si increases the surface hardness as well as surface N content as shown in Fig. 2, while comparing hardness at the same total N content, Si addition slightly reduces hardness than Fe–N martensite, while Al addition increases. At the same N content in solution, martensite hardness of Al- and Si- added alloys is higher, suggesting dispersion of AlN and α-Si3N4 contributes to the hardening of martensite. Additional martensite hardening by alloy carbide dispersion was also reported in carbide dispersion during carburizing in 12 Cr44) or 10Mo-0.5V,45) and hardness reaches about 1100 HV. Therefore, a similar heat treatment for the austenitic nitriding and quenching treatment will be possible. As in the case of carburizing, post-tempering is generally applied for austenitic nitriding.3) One benefit of using N martensite instead of C martensite is considerable softening resistance or even secondary hardening by the formation of M–N clusteing during tempering.46,47)
Microstructures and hardness in pure iron and Fe-1 mass%M (M = Mn, Cr, Al, Si, Mo) alloys austenitic-nitrided at 1123 K and subsequently quenched were investigated. The main conclusions are summarized as follows:
(1) Adding Al and Si promotes surface hardening, while the thickness of the austenite zone transformed into martensite is decreased. On the other hand, Mn addition increases both surface hardness and the thickness of the austenite zone. In contrast, Cr addition reduces hardness and thickness.
(2) AlN (wurtzite), α-Si3N4, and CrN (B1) precipitation occurs in the ferrite matrix and the austenite zone in Al-, Si- and Cr added alloys, respectively. A peculiar (austenite + α-Si3N4) lamellar structure develops in the Si-added alloy via isothermal eutectoid transformation. Those microstructure evolutions are consistent with the isothermal phase diagram.
(3) The N martensite hardness of pure iron and Mn-added alloy is similar at the same N content, indicating that the hardening by Mn addition resulted from an increment of N content due to Mn–N attractive interaction. Al and Si-added alloys exhibit higher martensite hardness than Fe–N martensite at the same N content in solution. The extra hardening of martensite by Al and Si might be caused by alloy nitrides, as previously reported in carbide-dispersion carburizing. Intragranular ferrite transformation on CrN during final quenching leads to softening by Cr addition.
This work was supported by the JST Collaborative Research Based on Industrial Demand (Grant No. JPMJSK1613, Japan), JST FOREST Program (Grant No. JPMJFR203W, Japan), MEXT Program Data Creation and Utilization Type Material Research and Development Project (Grant Number JPMXP1122684766). T.F. and G.M. gratefully acknowledge the support provided by the MEXT through Grants-in-Aid for Scientific Research (A) (No. 17H01330, 2017–2019), Grant-in-Aid for Scientific Research (B) (No. 19H02473, 2019–2021), Grant-in-Aid for Scientific Research on Innovative Areas (Research in a proposed research area) (No. 18H05456, 2018–2022), Grant-in-Aid for Challenging Research (Exploratory) (No. 21K18803, 2021–2022),. The Tohoku University Microstructural Characterization Platform in Nanotechnology Platform Project, sponsored by the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan is also acknowledged.
